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The influence of hot deformation on the microstructure of Ni-Al and Ni-Al-Fe alloys with small Ti2B addition
Journalof
ELSEVIER
Materials Processing Technology
Journal of Materials ProcessingTechnology60 (1996) 317-323
The influence of hot deformation on the microstructure of Ni-Al and Ni-A1-Fe alloys with small Ti2B addition. * T. Czeppe x , S. Szczepanik~ x Institute of Metallurgy and Materials Science, Polish Academy of Sciences, 25 Reymonta St., 30-059 Krak6w, Poland = University of Mining and Metallurgy, 30 Mickiewicza Av., 30-059 Krak6w, Poland
Abstract
Results of microstructure investigation of Ni-A1 and Ni-A1-Fe alloys with addition of 0.5at.% Ti2B after hot deformation are presented. Ti partially dissolves in the matrix, promoting formation of a discontinuous phase and precipitates of 7' phase of L12 structure during hot deformation of Ni-Al and Ni-A1-4.3at.%Fe alloys. In all investigated alloys precipitates of different compositions, enriched in Ti, were noticed. A compression test proved that due to the grain boundaries strengthening, the fracture mechanism changes into a transcrystalline one, improving fracture toughness of the hot deformed alloy. Fracture occurs in the deformation induced L 10 martensite.
Keywords: hot upsetting, microstructure, fracture, NiAl, NiA1Fe.
1. Introduction
The nonstoichiometric intermetallic NiA1 compound is thought to be promising as a future constructive material for high temperature application [1]. In the range of temperature which strongly depends on Ni contents, the 13 phase undergoes the thermoelastic martensitic transformation [2]. This makes the alloy potentially suitable as a shape memory material [1]. However, at the chemical composition far from stoichiometry, at temperature lower than 600K, NiAl exhibits lack of plasticity and very high brittleness which makes the application of this alloy, either as constructive or shape memory material, rather problematic. Many attempts to increase the poor ductility of NiA1 were undertaken by micro- and macro- alloying [3,4].Until now, the problem looks as if it has not been solved yet. Ti as a former of the brittle 7' phase does not improve the ductility of NiAl [5], in larger content Ti suppresses martensitic transformation, leading to the formation ofmetastable phases, not observed in the NiA1 alloy [6]. Also the studies of the influence of boron on the stoichiometric alloys showed that the plasticity was not better [4,7], althougt it altered the fracture mode. The studies of 2 wt.% additions of Ti2B suggested improvement of high temperature plasticity of NiAI alloys, but not at room temperature [8]. The present paper is aimed at the study of the influence of 0.5at% addition of Ti:B on the microstructure induced by hot
* This research was supported by the Polish State Commitee for Scientific Research under grant No 3P 40706306 0924-0136/96/$15.00 © 1996 Elsevier ScienceS.A. All fights reserved PI10924-0136 (96) 02348-5
upsetting in the Ni-A1 and Ni-Al-Fe alloys and on the fracture properties caused by microstructure of Ni-A1 alloy.
1
//12
•
//
\ \ \ \
\ \ \
Fig. 1. Schematic diagram of the deformation process: 1- sample, 110,h height before and after deformation, 2- carbon steel die.
318
1]. Czeppe, S. Szczepanik / Journal of Materials Processing Technology 60 (1996) 317-323
2. Experimental
3. Results
The alloys were prepared from high purity Ni, A1, Fe and A1-TizB master alloy by melting in a Balzers furnace under argon atmosphere and casting into a chill mould in vacuum. Samples were hot upset in a "soft" carbon steel die at temperature 1273 K and quickly cooled. Schematic diagram of deformation process is presented in Fig. 1. The deformation was defined as: ~ = ((h - h0)/ho)xl00% where ho and h are heights of the sample before and after the deformation, respectively. Various degrees of deformation were applied to the samples.The NiA1 alloy in "as cast" state of dendritic microstructure, was deformed by 20 and 40 %. The same alloy, after homogenisation at 1473 K in the 13 phase range was deformed by 25 and 50%. All Ni-A1-Fe alloys were deformed in the 13phase, after homogenisation at 1473 K, by 25%. The microstructure of alloys was observed using transmission electron microscopes Philips 301 (100kV) and CM20 (200 kV). The microanalysis of thin foils was performed with EDS Exl Link system while the microcomposition of the bulk was analy- sed with EDS system of Philips XL30 SEM. The compression test was performed with the Instron-6025.The chemical composition of alloys was carefully controlled after homogenisation and after plastic deformation. Results are presented in Table 1. As the boron content could not be measured, the nominal addition of 0.2 at.% was assumed. The error in the Ti content determination was estimated to be lower than 0.1 at.%. The impurities were below the level of detection by the EDS method.
3.1. The microstructure o f the alloys after hot deformation. 3.1.1. Ni- A1 (A1) alloy. The optical microstructure of the sample deformed in "as cast" state is presented in Fig.2a. The hot deformation produces small grains which are elongated perpendicular to the compression axis with a slightly different average grain diameter, depending on the distance from the central part of the sample. The compositional nonhomogenity caused by the initial dendritic structure remined after the deformation. In Fig.2a precipitates are also visible. After 5 h ofhomogenisation at 1473 K the average grain size was about 300 ~tm, the material was fully recrystallised and homogenised with the thermally stable, Ti-rich precipitates. This material was further used as a standard, for comparison with the samples deformed in the 13 phase. In spite of high brittleness of the 13 phase, the samples were successfully deformed by 25 and 50%. After 25% deformation, the material revealed small grains of average size of 80 ~na Table 1. Chemical composition in at.% and designation of alloys, sumed value)
Al B1 B2 B3
A1
Fe
Ti
B
Ni
34.8 36.5 35.1 35.9
4.3 10.7 19.1
0.2 0.1 0.1 0.1
0.2* 0.2* 0.2* 0.2*
bal. bal. bal. bal.
-
as-
Fig. 2. Microstructure of the A1 alloy after hot deformation, a) "as cast sample", 40% of deformation, optical microscope, x125, b) sample deformed in the 13 phase by 25%, optical microscope, xl00, c) sample deformed in the 13 phase by 25%, SEM, x2000, dark precipitates are enriched in Ti, ,/' phase forms light precipitates.
T. Czeppe, S. Szczepanik / Journal of Materials Processing Technology 60 (1996) 317-323
measured in the direction of the compression axis and about 30% more in the perpendicular plane (Fig. 2b,c). The discontinuous phase at grain boundaries and a new type of precipitates in the grains, formed during hot deformation process are clearly seen in Fig.2c. The sample after 50% deformation exhibited large grains, strongly elongated in the directions perpendicular to the compression axis, of average size about 10 times larger than in the previous sample (Fig 3). Small subgrains at the recrystallised grain boundaries and transgranular microcracks, formed during hot deformation are also visible. The SEM observation revealed also small precipitates in the grains (Fig.3) and phase at the grain boundaries. The examination of the matrix structure by the TEM showed the B2 structure of long range atomic order. In the case of sample deformed by 25%, small subgrains together with a pronounced dislocation structure in the grains were observed (Fig.4a,b). Using EDS method three compositions of the oval-shaped, Tirich precipitates were determined in at.%: 1). Ti - 5Ni - 1A1, 2).
319
Ti - 25 Ni - 10A1, 3). Ti - 40Ni - 15A1 (Fig.6a). The first of these compositions refers to the Ti2B particles. The L12 structure of a discontinuous phase or individual precipitates at the grain boundaries (Fig.5), and of the considerable amount of lenticular precipitates (Fig.6b) in the matrix of the sample deformed by 25%, was identified as the structure of NisAlxTi~.x(7') phase. The composition of the discontinuous phase was found to be in the range of (68-70)Ni - (29 - 31)A1 - (0.2-0.7)Ti (at.%), containing less Ni than the equilibrium 7' phase [9 ]; the composition of the lenticular precipitates was Ni - 24A1 - 0.5Ti (at%), typical for the 7' phase. In the sample deformed by 50%, composition of precipitates was in the same range as this estabilished for the discontinuous phase. 3.1.2. Ni-AI-Fe, B1 - B3 alloys.
After 25% of deformation, alloys B2 and B3 presented similar microstructure with small grains, 80 ~xn of average diameter, like in the A1 alloy . However, the difference in grains size in
Fig. 3. Microstructure of the A1 alloy, hot deformed in the 15phase by 50%, a) optical microscope, x 60, b) SEM, x400, light points are 7' precipitates, dark points - precipitates enriched in Ti.
Fig. 4. a) Electron microstructure of the A1 alloy after 25% hot deformation, b) diffraction pattern of the B2 structure, foil orientation [001].
320
T. Czeppe, S. Szczepanik / Journal of Materials Processing Technology 60 (1996) 317-323
Fig. 5. Phase at the grain boundaries. Alloy A1, sample hot deformed by 25%, a) scanning electron microscope, x5000, b) transmission electron microscope. the direction of compression was less pronounced (Fig.7). The microstructure of alloy B 1 was different. In that case, dynamic recrystallisation took place during deformation, leaving small grains among the recrystaUised ones (Fig.,8). Except the Ti-rich precipitates which were found in all samples, a different composition of precipitates was observed in the B1 alloy. It was Ni 11A1 - 7Ti - 6Fe (at.%) which was characteristic for '{' phase, like in the case of the A1 alloy, while in two other alloys the precipitates contain nearly the same content of Ni and Fe like in the matrix phase but are from 10 to 100 times richer in Ti. The average composition of that kind of precipitates in alloy B2 and
B3 was Ni -25.8A1 - 11.5Fe- 7.2Ti and Ni -34.2A1 - 19.0Fe 1.9Ti (at.%), respectively. 3.2. The effect o f microstructure on the fracture and microhardness o f alloys. To elucidate the influence of different microstructures produced during hot deformation, two samples of alloy A1, homogenised and 25% deformed, underwent the compression test at the compression rate 0.05 mm/min, at the room temperature. The results, presented in Fig.9 (elastic part of strain in the
B
t~:
Fig, 6. Transmission electron microstructure of the precipitates in A1 alloy, hot deformed by 25% in the 13 phase, a) Ti - rich precipitate, b) y' precipitate, SADP, foil orientation [001] L12 .
T. Czeppe, S. Szczepanik / Journal of Materials Processing Technology 60 (1996) 317-323
321
°
lv, /
x,
t %
J
t
Fig. 7. Optical microstructure of B2 alloy after hot deformation by 25%, x200.
Fig. 8. Electron microstructure of B1 alloy after 25% of hot deformation, small grains at the boundary of two recrystallised ones.
figure was not subtracted) showed that in the case of the sample hot deformed by 25%, fracture stress was 1800 MPa, three times higher than for the same material after homogenisation. A small range of plastic deformation was also visible. Optical and SEM observations, made on prepolished surfaces of the deformed samples showed difference in the fracture mechanism. The fracture was generally intergranular in the sample homogenised (Fig.10a) when in the sample just after the hot deformation - transgranular (Fig.10b). Careful examination of the grain boundaries in the latter sample, proved that cracks did not passed through the discontinuous phase at the grain boundaries, and that fracture was noticed at the grain boundaries only at the places where the phase or precipitates did not form.
2000
I
~
t
,
1
II - sample with precipitates
1600 --
n
t I ~-
1200
I/)
(/') 800
I - sample
homogenised
400
0
r
),00
I
'
0.40
I 0.80
Strain
Fig. 9 Stress - strain to fracture curves, (~=Ah/h 0 ).
Fig. 10. SEM microstructure of the A1 alloy after fracture, a) sample homogenised, b) sample with precipitates.
322
T. Czeppe, S. Szczepanik / Journal of Materials Processing Technology 60 (1996) 317-323
The compression - fracture surfaces are presented in Fig. 11. In both samples fracture surfaces are not smooth but reveal sort of lamellar structure. The TEM examination of the sample structures after deformation proved that the fracture occurred in the deformation induced martensite, predominantly of the L 10 structure where the fracture mechanism must have incorporated in-
ternal twins and intervariant boundaries of the martensite. The comparison of microhardness of samples (Tab.2) confirmed that sample of the A1 alloy, with the precipitates of 7' phase, was less hard than the homogenised one, when microhardness of the B1 alloy was in the same range, but the microhardness of alloys 132 and 133 was lower.
Table 2. HV10 microhardness of investigated alloys. Alloy Sample deformed and homogenised in the [3 phase 25 % deformation in the [3 phase
A1 413 390
B1
401
B2
348
B3
354
4.Conclusions
Fig. 11. Compression - fracture surfaces, alloy A1, SEM, x 550 a) sample homogenised, b) sample with precipitates.
1. The applied method of high temperature deformation at 1273K leads to the formation of the microstructure consisting of small (80 lam average diameter) elongated grains in the Ni-A1 and Ni-A1-Fe alloys. During the hot deformation a dynamic recrystallisation took place. A higher degree of recrystallisation was observed in the Ni-A1 alloy deformed by 50% and in the Ni-A1-4.3at.%Fe one. 2. The addition of 0.5at.% of TizB to Ni-A1 and Ni-A1-Fe alloys leads to dissolution of about 2/3 of Ti in the matrix. The rest of Ti enters precipitates enriched in Ti or Ni, of different compositions and morphology. The composition of the Ti-enriched precipitates in the case of alloys with additions of 10.7at.% Fe and 19.1at.% Fe suggests that Ti mainly replaces the A1 atoms. Boron content could not be detected but it should be assumed that it locates preferably at the grain boundaries and in the precipitates [7 ]. 3. Ti, as the 7' phase former [5] promoted formation of the discontinuous phase or precipitates, located at the grain boundaries in the Ni-A1 samples hot deformed in the 13 phase. This phase contains slightly less Ni than the equilibrium 7' phase but its structure is also L12 .Precipitates of similiar composition were found in Ni-A1-4.1at.%Fe alloy. The precipitates having the 7' phase composition and structure, also formed in the grains of Ni-A1 samples, deformed in the 13phase. 4. The compression - fracture experiment proved that the fracture toughness of the Ni-A1 sample containing the discontinuous phase at the grain boundaries and the 7' precipitates inside the grains, was three times higher than that of the sample homogenised after the hot deformation. Two different fracture modes were observed: transgranular in the former and intergranular in the latter case. The change of the fracture mode is due to the presence of the discontinuous phase at the grain bounadries. The role of the boron addition is not clear in that case, as it should change the fracture mode also in the case of the homogenised sample, at room temperature [4, 7]. Much smaller grain size in the sample non homogenised after hot deformation may also add to higer fracture toughness, as was presented for stoichiometric NiA1 [10]. 5. It was proved by the fracture surfaces and TEM observations, that fracture occured in both samples inside the martensite in-
T. Czeppe, S. Szczepanik/ Journal of Materials Processing Technology 60 (1996) 317-323
duced by deformation, predominantly of the L10 structure. The fracture mechanism must have incorporated internal twins and intervariant boundaries in the martensitic phase. 6. The microhardness measurements showed that at room ttemperature the Ni-A1 sample after homogenisation was harder than after the hot deformation, whereas Ni-Al-10.7at.%Fe and Ni-AI-19.1 at.%Fe alloys were less hard.
Acknowledgement Authors wish to thank Dr. H. Paul for the performing of the compression test and valuable discussions.
References [1] D.B. Miracle, ActaMetalL Mater., 41 (1993) 649
323
[2] S. Chakravorty and C.M. Wayman, Metall. Trans., 7A (1976) 555 [3] K. Ishida, R. Kainuma, N. Ueno and T. Nishizawa, Metall. Trans., 22A (1991) 441 [4] E.P. George and C.T. Liu, J. Mater. Res., 5 (1990) 754 [5] C.E. Jia, R. Ishida and T. Nishizawa, Metall. andMater. Trans. A, 25A (1994) 473 [6] R. Yang, S.V. Parker, J.A. Leake and, R.W. Cahn, Alloy Modeling and Design, (eds).G.M.Stocks, P.E.A. Turchi, The Minerals,Metals & Materials Society, 1994, p303 [7] R. Jayaram and M.K. Miller, Acta metall.mater. 42 (1994) 1561 [8] Lee In-Gyu, A.K. Gosha, R. Ray, and S. Jha, Metall. and Mater. Trans.A 25A (1994) 2017 [9] G. Petzow and G. Effenberg (eds.), Ternary Alloys, 1993, vol.5, p.317, vol.8, p.20 [10], P. Nagpal and I. Baker, ScriptaMetall. 24 (1990) 2381
ELSEVIER
Materials Processing Technology
Journal of Materials ProcessingTechnology60 (1996) 317-323
The influence of hot deformation on the microstructure of Ni-Al and Ni-A1-Fe alloys with small Ti2B addition. * T. Czeppe x , S. Szczepanik~ x Institute of Metallurgy and Materials Science, Polish Academy of Sciences, 25 Reymonta St., 30-059 Krak6w, Poland = University of Mining and Metallurgy, 30 Mickiewicza Av., 30-059 Krak6w, Poland
Abstract
Results of microstructure investigation of Ni-A1 and Ni-A1-Fe alloys with addition of 0.5at.% Ti2B after hot deformation are presented. Ti partially dissolves in the matrix, promoting formation of a discontinuous phase and precipitates of 7' phase of L12 structure during hot deformation of Ni-Al and Ni-A1-4.3at.%Fe alloys. In all investigated alloys precipitates of different compositions, enriched in Ti, were noticed. A compression test proved that due to the grain boundaries strengthening, the fracture mechanism changes into a transcrystalline one, improving fracture toughness of the hot deformed alloy. Fracture occurs in the deformation induced L 10 martensite.
Keywords: hot upsetting, microstructure, fracture, NiAl, NiA1Fe.
1. Introduction
The nonstoichiometric intermetallic NiA1 compound is thought to be promising as a future constructive material for high temperature application [1]. In the range of temperature which strongly depends on Ni contents, the 13 phase undergoes the thermoelastic martensitic transformation [2]. This makes the alloy potentially suitable as a shape memory material [1]. However, at the chemical composition far from stoichiometry, at temperature lower than 600K, NiAl exhibits lack of plasticity and very high brittleness which makes the application of this alloy, either as constructive or shape memory material, rather problematic. Many attempts to increase the poor ductility of NiA1 were undertaken by micro- and macro- alloying [3,4].Until now, the problem looks as if it has not been solved yet. Ti as a former of the brittle 7' phase does not improve the ductility of NiAl [5], in larger content Ti suppresses martensitic transformation, leading to the formation ofmetastable phases, not observed in the NiA1 alloy [6]. Also the studies of the influence of boron on the stoichiometric alloys showed that the plasticity was not better [4,7], althougt it altered the fracture mode. The studies of 2 wt.% additions of Ti2B suggested improvement of high temperature plasticity of NiAI alloys, but not at room temperature [8]. The present paper is aimed at the study of the influence of 0.5at% addition of Ti:B on the microstructure induced by hot
* This research was supported by the Polish State Commitee for Scientific Research under grant No 3P 40706306 0924-0136/96/$15.00 © 1996 Elsevier ScienceS.A. All fights reserved PI10924-0136 (96) 02348-5
upsetting in the Ni-A1 and Ni-Al-Fe alloys and on the fracture properties caused by microstructure of Ni-A1 alloy.
1
//12
•
//
\ \ \ \
\ \ \
Fig. 1. Schematic diagram of the deformation process: 1- sample, 110,h height before and after deformation, 2- carbon steel die.
318
1]. Czeppe, S. Szczepanik / Journal of Materials Processing Technology 60 (1996) 317-323
2. Experimental
3. Results
The alloys were prepared from high purity Ni, A1, Fe and A1-TizB master alloy by melting in a Balzers furnace under argon atmosphere and casting into a chill mould in vacuum. Samples were hot upset in a "soft" carbon steel die at temperature 1273 K and quickly cooled. Schematic diagram of deformation process is presented in Fig. 1. The deformation was defined as: ~ = ((h - h0)/ho)xl00% where ho and h are heights of the sample before and after the deformation, respectively. Various degrees of deformation were applied to the samples.The NiA1 alloy in "as cast" state of dendritic microstructure, was deformed by 20 and 40 %. The same alloy, after homogenisation at 1473 K in the 13 phase range was deformed by 25 and 50%. All Ni-A1-Fe alloys were deformed in the 13phase, after homogenisation at 1473 K, by 25%. The microstructure of alloys was observed using transmission electron microscopes Philips 301 (100kV) and CM20 (200 kV). The microanalysis of thin foils was performed with EDS Exl Link system while the microcomposition of the bulk was analy- sed with EDS system of Philips XL30 SEM. The compression test was performed with the Instron-6025.The chemical composition of alloys was carefully controlled after homogenisation and after plastic deformation. Results are presented in Table 1. As the boron content could not be measured, the nominal addition of 0.2 at.% was assumed. The error in the Ti content determination was estimated to be lower than 0.1 at.%. The impurities were below the level of detection by the EDS method.
3.1. The microstructure o f the alloys after hot deformation. 3.1.1. Ni- A1 (A1) alloy. The optical microstructure of the sample deformed in "as cast" state is presented in Fig.2a. The hot deformation produces small grains which are elongated perpendicular to the compression axis with a slightly different average grain diameter, depending on the distance from the central part of the sample. The compositional nonhomogenity caused by the initial dendritic structure remined after the deformation. In Fig.2a precipitates are also visible. After 5 h ofhomogenisation at 1473 K the average grain size was about 300 ~tm, the material was fully recrystallised and homogenised with the thermally stable, Ti-rich precipitates. This material was further used as a standard, for comparison with the samples deformed in the 13 phase. In spite of high brittleness of the 13 phase, the samples were successfully deformed by 25 and 50%. After 25% deformation, the material revealed small grains of average size of 80 ~na Table 1. Chemical composition in at.% and designation of alloys, sumed value)
Al B1 B2 B3
A1
Fe
Ti
B
Ni
34.8 36.5 35.1 35.9
4.3 10.7 19.1
0.2 0.1 0.1 0.1
0.2* 0.2* 0.2* 0.2*
bal. bal. bal. bal.
-
as-
Fig. 2. Microstructure of the A1 alloy after hot deformation, a) "as cast sample", 40% of deformation, optical microscope, x125, b) sample deformed in the 13 phase by 25%, optical microscope, xl00, c) sample deformed in the 13 phase by 25%, SEM, x2000, dark precipitates are enriched in Ti, ,/' phase forms light precipitates.
T. Czeppe, S. Szczepanik / Journal of Materials Processing Technology 60 (1996) 317-323
measured in the direction of the compression axis and about 30% more in the perpendicular plane (Fig. 2b,c). The discontinuous phase at grain boundaries and a new type of precipitates in the grains, formed during hot deformation process are clearly seen in Fig.2c. The sample after 50% deformation exhibited large grains, strongly elongated in the directions perpendicular to the compression axis, of average size about 10 times larger than in the previous sample (Fig 3). Small subgrains at the recrystallised grain boundaries and transgranular microcracks, formed during hot deformation are also visible. The SEM observation revealed also small precipitates in the grains (Fig.3) and phase at the grain boundaries. The examination of the matrix structure by the TEM showed the B2 structure of long range atomic order. In the case of sample deformed by 25%, small subgrains together with a pronounced dislocation structure in the grains were observed (Fig.4a,b). Using EDS method three compositions of the oval-shaped, Tirich precipitates were determined in at.%: 1). Ti - 5Ni - 1A1, 2).
319
Ti - 25 Ni - 10A1, 3). Ti - 40Ni - 15A1 (Fig.6a). The first of these compositions refers to the Ti2B particles. The L12 structure of a discontinuous phase or individual precipitates at the grain boundaries (Fig.5), and of the considerable amount of lenticular precipitates (Fig.6b) in the matrix of the sample deformed by 25%, was identified as the structure of NisAlxTi~.x(7') phase. The composition of the discontinuous phase was found to be in the range of (68-70)Ni - (29 - 31)A1 - (0.2-0.7)Ti (at.%), containing less Ni than the equilibrium 7' phase [9 ]; the composition of the lenticular precipitates was Ni - 24A1 - 0.5Ti (at%), typical for the 7' phase. In the sample deformed by 50%, composition of precipitates was in the same range as this estabilished for the discontinuous phase. 3.1.2. Ni-AI-Fe, B1 - B3 alloys.
After 25% of deformation, alloys B2 and B3 presented similar microstructure with small grains, 80 ~xn of average diameter, like in the A1 alloy . However, the difference in grains size in
Fig. 3. Microstructure of the A1 alloy, hot deformed in the 15phase by 50%, a) optical microscope, x 60, b) SEM, x400, light points are 7' precipitates, dark points - precipitates enriched in Ti.
Fig. 4. a) Electron microstructure of the A1 alloy after 25% hot deformation, b) diffraction pattern of the B2 structure, foil orientation [001].
320
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Fig. 5. Phase at the grain boundaries. Alloy A1, sample hot deformed by 25%, a) scanning electron microscope, x5000, b) transmission electron microscope. the direction of compression was less pronounced (Fig.7). The microstructure of alloy B 1 was different. In that case, dynamic recrystallisation took place during deformation, leaving small grains among the recrystaUised ones (Fig.,8). Except the Ti-rich precipitates which were found in all samples, a different composition of precipitates was observed in the B1 alloy. It was Ni 11A1 - 7Ti - 6Fe (at.%) which was characteristic for '{' phase, like in the case of the A1 alloy, while in two other alloys the precipitates contain nearly the same content of Ni and Fe like in the matrix phase but are from 10 to 100 times richer in Ti. The average composition of that kind of precipitates in alloy B2 and
B3 was Ni -25.8A1 - 11.5Fe- 7.2Ti and Ni -34.2A1 - 19.0Fe 1.9Ti (at.%), respectively. 3.2. The effect o f microstructure on the fracture and microhardness o f alloys. To elucidate the influence of different microstructures produced during hot deformation, two samples of alloy A1, homogenised and 25% deformed, underwent the compression test at the compression rate 0.05 mm/min, at the room temperature. The results, presented in Fig.9 (elastic part of strain in the
B
t~:
Fig, 6. Transmission electron microstructure of the precipitates in A1 alloy, hot deformed by 25% in the 13 phase, a) Ti - rich precipitate, b) y' precipitate, SADP, foil orientation [001] L12 .
T. Czeppe, S. Szczepanik / Journal of Materials Processing Technology 60 (1996) 317-323
321
°
lv, /
x,
t %
J
t
Fig. 7. Optical microstructure of B2 alloy after hot deformation by 25%, x200.
Fig. 8. Electron microstructure of B1 alloy after 25% of hot deformation, small grains at the boundary of two recrystallised ones.
figure was not subtracted) showed that in the case of the sample hot deformed by 25%, fracture stress was 1800 MPa, three times higher than for the same material after homogenisation. A small range of plastic deformation was also visible. Optical and SEM observations, made on prepolished surfaces of the deformed samples showed difference in the fracture mechanism. The fracture was generally intergranular in the sample homogenised (Fig.10a) when in the sample just after the hot deformation - transgranular (Fig.10b). Careful examination of the grain boundaries in the latter sample, proved that cracks did not passed through the discontinuous phase at the grain boundaries, and that fracture was noticed at the grain boundaries only at the places where the phase or precipitates did not form.
2000
I
~
t
,
1
II - sample with precipitates
1600 --
n
t I ~-
1200
I/)
(/') 800
I - sample
homogenised
400
0
r
),00
I
'
0.40
I 0.80
Strain
Fig. 9 Stress - strain to fracture curves, (~=Ah/h 0 ).
Fig. 10. SEM microstructure of the A1 alloy after fracture, a) sample homogenised, b) sample with precipitates.
322
T. Czeppe, S. Szczepanik / Journal of Materials Processing Technology 60 (1996) 317-323
The compression - fracture surfaces are presented in Fig. 11. In both samples fracture surfaces are not smooth but reveal sort of lamellar structure. The TEM examination of the sample structures after deformation proved that the fracture occurred in the deformation induced martensite, predominantly of the L 10 structure where the fracture mechanism must have incorporated in-
ternal twins and intervariant boundaries of the martensite. The comparison of microhardness of samples (Tab.2) confirmed that sample of the A1 alloy, with the precipitates of 7' phase, was less hard than the homogenised one, when microhardness of the B1 alloy was in the same range, but the microhardness of alloys 132 and 133 was lower.
Table 2. HV10 microhardness of investigated alloys. Alloy Sample deformed and homogenised in the [3 phase 25 % deformation in the [3 phase
A1 413 390
B1
401
B2
348
B3
354
4.Conclusions
Fig. 11. Compression - fracture surfaces, alloy A1, SEM, x 550 a) sample homogenised, b) sample with precipitates.
1. The applied method of high temperature deformation at 1273K leads to the formation of the microstructure consisting of small (80 lam average diameter) elongated grains in the Ni-A1 and Ni-A1-Fe alloys. During the hot deformation a dynamic recrystallisation took place. A higher degree of recrystallisation was observed in the Ni-A1 alloy deformed by 50% and in the Ni-A1-4.3at.%Fe one. 2. The addition of 0.5at.% of TizB to Ni-A1 and Ni-A1-Fe alloys leads to dissolution of about 2/3 of Ti in the matrix. The rest of Ti enters precipitates enriched in Ti or Ni, of different compositions and morphology. The composition of the Ti-enriched precipitates in the case of alloys with additions of 10.7at.% Fe and 19.1at.% Fe suggests that Ti mainly replaces the A1 atoms. Boron content could not be detected but it should be assumed that it locates preferably at the grain boundaries and in the precipitates [7 ]. 3. Ti, as the 7' phase former [5] promoted formation of the discontinuous phase or precipitates, located at the grain boundaries in the Ni-A1 samples hot deformed in the 13 phase. This phase contains slightly less Ni than the equilibrium 7' phase but its structure is also L12 .Precipitates of similiar composition were found in Ni-A1-4.1at.%Fe alloy. The precipitates having the 7' phase composition and structure, also formed in the grains of Ni-A1 samples, deformed in the 13phase. 4. The compression - fracture experiment proved that the fracture toughness of the Ni-A1 sample containing the discontinuous phase at the grain boundaries and the 7' precipitates inside the grains, was three times higher than that of the sample homogenised after the hot deformation. Two different fracture modes were observed: transgranular in the former and intergranular in the latter case. The change of the fracture mode is due to the presence of the discontinuous phase at the grain bounadries. The role of the boron addition is not clear in that case, as it should change the fracture mode also in the case of the homogenised sample, at room temperature [4, 7]. Much smaller grain size in the sample non homogenised after hot deformation may also add to higer fracture toughness, as was presented for stoichiometric NiA1 [10]. 5. It was proved by the fracture surfaces and TEM observations, that fracture occured in both samples inside the martensite in-
T. Czeppe, S. Szczepanik/ Journal of Materials Processing Technology 60 (1996) 317-323
duced by deformation, predominantly of the L10 structure. The fracture mechanism must have incorporated internal twins and intervariant boundaries in the martensitic phase. 6. The microhardness measurements showed that at room ttemperature the Ni-A1 sample after homogenisation was harder than after the hot deformation, whereas Ni-Al-10.7at.%Fe and Ni-AI-19.1 at.%Fe alloys were less hard.
Acknowledgement Authors wish to thank Dr. H. Paul for the performing of the compression test and valuable discussions.
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